Properties and Corrosion 349 tests indicated the presence of stress-enhanced oxidation at 1000°C, with failure times ranging from 19 to 93 hr at an applied load of 138 MPa, and from 14 to 31 hr at an applied load of 276 MPa. Losses in strength at temperatures greater than 1200°C were attributed to the softening of the glassy grain boundary phase, which leads to creep by grain boundary sliding. Samples exposed to oxidation at 1200°C at an applied load of 344 MPa, did not fail, even after 260 hr, although some slight deformation had occurred. In an effort to determine the effects of oxidation upon the flexural strength of Si 3 N 4 , Kim and Moorhead [8.30] evaluated the room-temperature four-point bend strength of HIP-SN (with 6 wt.% Y 2 O 3 and 1.5 wt.% Al 2 O 3 ) after exposure in either H 2 /H 2 O or Ar/O 2 at 1400°C for 10 hr. In both atmospheres, the strength was dependent on the amount of oxidant present. However, the actual variation in strength was different, depending upon the alteration of the surface layers formed and their characteristics. In the H 2 /H 2 O atmosphere at low pH 2 O, a nonprotective and not well-attached glass-like layer containing crystalline Y 2 Si 2 O 7 formed. Because this layer was relatively uniform with no new strength-limiting flaws being formed (although some large bubbles were found at the surface/substrate interface), the maximum reduction in strength was limited to about 20% at a pH 2 O of 2×10 -5 MPa. A significant strength increase occurred as the pH 2 O was increased, which the authors attributed to blunting of preexisting cracks by the interfacial silicate phase. This silicate phase was a continuous dense layer of Y 2 Si 2 O 7 containing small isolated bubbles believed to be formed by nitrogen generation during oxidation of the Si 3 N 4 . In the Ar/O 2 atmosphere, a similar reduction and subsequent increase in strength was not found. Instead, at low pO 2 , an increase in strength occurred with increasing pO 2 . The maximum strength occurred at pO 2 (10 -5 MPa) that yielded the greatest weight loss. Even at low pO 2 , a surface reaction product of Y 2 Si 2 O 7 formed in isolated pockets at grain junctions, presumably by the reaction of Y 2 O 3 solid with SiO gas. Kim and Moorhead attributed the increased Copyright © 2004 by Marcel Dekker, Inc. 350 Chapter 8 strengths observed to the formation of more Y 2 Si 2 O 7 as the pO 2 increased. At approximately a pO 2 of 10 -5 MPa, where the maximum strength was observed, the Y 2 Si 2 O 7 layer became interconnected and, although not continuous, blunted strength limiting flaws. At higher pO 2 , where weight gains were observed and a continuous layer containing Y 2 Si 2 O 7 and cristobalite formed, the increase in strength was not as significant. In this region, competition between crack blunting and formation of new flaws (cracks and bubbles) was suggested as the reason for the slightly lower strengths. This particular study by Kim and Moorhead pointed out very well the effects that the surface layer characteristics have upon the mechanical properties. Similar strength increases were found by Wang et al. [8.31] for two silicon nitride materials, one containing 13.9% Y 2 O 3 plus 4.5% Al 2 O 3 and the other containing 15% Y 2 O 3 plus 5% Al 2 O 3 , when exposed to air at 1200°C for 1000 hr prior to strength testing at 1300°C. Strength increases as high as 87% were reported when compared to the unoxidized 1300°C strength, although the preoxidized 1300°C strength was slightly less than the unoxidized room temperature strength. Wang et al. attributed these strength increases to healing of surface flaws and crack blunting during oxidation, along with purification of the grain boundaries that raised the viscosity of the glassy boundary phase. These beneficial effects were not present when oxidation was conducted at 900°C. Lange and Davis [8.32] have suggested that oxidation can lead to surface compressive stresses that, if optimum, may lead to increased apparent strengths. If the compressive stresses become too severe, then spalling may occur leading to lowered strengths. They demonstrated this concept with Si 3 N 4 doped with 15% and 20% CeO 2 exposed to oxidation in air, at temperatures ranging from 400 to 900°C. The apparent critical stress intensity factor (K a ) increased for short exposure times at 400, 500, and 600°C. This increase in K a was attributed to oxidation of the Ce-apatite secondary phase and subsequent development of a surface compressive layer. At longer times ( ~ 8 hr) and the two higher temperatures, surface spalling caused a decrease in K a . At higher Copyright © 2004 by Marcel Dekker, Inc. Properties and Corrosion 351 temperatures (i.e., 1000°C), the compressive stresses that may cause spalling were relieved by extrusion of the oxide product from the interior of the material. Thus, prolonged oxidation at 1000°C did not degrade this material. Oxynitrides In a study of β’ and O’ SiAlON solutions, O’Brien et al. [8.33] found that the oxygen (or nitrogen) content significantly affected the performance of these materials. The grain boundary glassy phase viscosity increased as the nitrogen content increased, which subsequently slowed the healing of flaws (see Chapter 2, Section 2.2.3 on Glasses and Chapter 6, Section 6.2 upon Silicate Glasses for a discussion of the effects of nitrogen upon durability). The higher viscosity glassy phase also trapped evolving gases more easily, creating additional flaws. In general, the mean retained flexural strengths after oxidation at 1273 K for 24 hr of the SiAlON solutions was higher than that of several silicon nitrides, with the strengths being generally proportional to the oxidation resistance. O’Brien et al. concluded that the retained strengths after oxidation were dependent upon the characteristics of the surface oxide layer that formed. At higher temperatures, the potential for flaw healing was dependent upon the amount and composition of the glassy phase formed. A zirconium oxynitride with the stoichiometry ZrO 2–2x N 4x/3 was reported by Claussen et al. [8.34] to form as a secondary phase in hot-pressed ZrO 2 –Si 3 N 4 . This phase readily oxidized to monoclinic ZrO 2 at temperatures greater than 500°C. Lange [8.35] used the volume change (about 4–5%) associated with this oxidation to evaluate the formation of a surface compression layer on silicon nitride compositions containing 5–30 vol.% zirconia. To develop the correct stress distribution for formation of the surface compressive layer, the secondary phase that oxidizes must be uniformly distributed throughout the matrix. When oxidized at 700°C for 5 hr, a material containing 20 vol.% ZrO 2 exhibited an increase in strength from 683 to 862 MPa. Lange attributed this increase in strength Copyright © 2004 by Marcel Dekker, Inc. 352 Chapter 8 to the oxidation-induced phase change of the zirconium oxynitride to monoclinic zirconia. 8.3.2 Degradation by Moisture Lifetimes that are predicted from different fatigue tests will vary. Slow crack growth has been reported by Kawakubo and Komeya [8.36] to accelerate under cyclic conditions, especially of the tension—compression type cycle at room temperature for sintered silicon nitride. They also reported a plateau at about 70–90% of the stress intensity factor, when crack velocity was plotted vs. K I . Three regions in the data were observed, very similar to that reported for glasses as shown in Fig. 8.1. As the materials studied had a glassy grain boundary phase, the fatigue mechanism was assumed to be the same as that reported for glassy materials [8.13] (i.e., stress corrosion cracking due to moisture in the air). Fett et al. [8.37] reported that at 1200°C, the lifetimes for cyclic loads were higher than for static loads. Tajima et al. [8.38] reported that a gas pressure sintered silicon nitride was resistant to slow crack growth up to 900°C, but then was susceptible to slow crack growth at 1000°C because of the softening of the glassy grain boundary phase. A higher fatigue resistance was reported for higher frequencies of the load cycle due to the viscoelastic nature of the glassy grain boundary phase. 8.3.3 Degradation by Other Atmospheres Carbides and Nitrides Clark [8.39] reported that Nicalon™ SiC fibers when aged in nitrogen or humid air at 1200°C for 2 hr, lost about one-half of their tensile strength. A more gradual strength decrease was observed for fibers that were exposed to hot argon. Although the time dependence of strength loss for the different aging environments was similar, the mechanisms causing strength loss were quite different. For exposure to nitrogen, Clark attributed the strength loss to crack propagation from existing Copyright © 2004 by Marcel Dekker, Inc. Properties and Corrosion 353 flaws; for exposure to argon, he attributed the loss to grain growth and porosity; and for exposure to humid air, he attributed the strength loss to fiber coalescence at the silica surface, to poor adherence of the surface silica layer, to a cracked crystalline silica surface layer, and to bubbles at the silica/fiber interface. Clark also pointed out that thermal stability should not be based solely upon weight change data, because for this fiber, the weight gain produced by oxidation to silica was offset by weight loss due to CO evolution. Siliconized, boron-doped, and aluminum-doped SiC samples were exposed to gaseous environments containing mixtures of predominantly N 2 , H 2 , and CO, representative of metallurgical heat-treatment atmospheres at 1300°C for up to 1000 hr by Butt et al. [8.40]. They reported significant strength losses for all three materials for times less than 100 hr when exposed to a gas mixture containing about 40% nitrogen. At longer exposure times, no additional strength loss occurred. The aluminum-doped SiC, unlike the other two, exhibited a slight strength increase after 1000 hr when exposed to a gas mixture containing 98.2% nitrogen. The strength losses were attributed primarily to pitting that was related to the presence of transition metal impurities. It has been shown by Li and Langley [8.41] that ceramic fibers composed of Si–C–N–O experienced various degrees of strength degradation when aged in atmospheres of various hot gases. The rate of strength loss experienced by fibers aged in these hot gases was related to the rate of diffusion of the gases formed by decomposition. The gases of decomposition (N 2 , CO, and SiO) diffused through the fiber porosity and any surface boundary layers present. The diffusion of these product gases can be controlled by aging the fibers in atmospheres of these gases. Thus, greater strength loss was exhibited when fibers were aged in argon compared to aging in nitrogen. This effect can be seen by examining the data of Table 8.2. Zirconia-Containing Materials Brinkman et al. [8.42] studied the effects of a diesel engine environment upon the strength of two commercial zirconias Copyright © 2004 by Marcel Dekker, Inc. Properties and Corrosion 355 samples after most of the reaction products were removed. Those samples for which the reaction products were not removed prior to strength testing exhibited no significant loss of strength, although an increase in scatter of the data was reported. Surface or corrosion pits were identified as the fracture origin for both types of SiC. In addition, the α-SiC exhibited grain boundary attack, whereas the siliconized-SiC exhibited oxidation of the silicon matrix and attack of the large SiC grains. In a study of the effects of molten salt upon the mechanical properties of silicon nitride, Bourne and Tressler [8.44] reported that hot-pressed silicon nitride exhibited a more severe degradation in flexural fracture strength than did reaction sintered silicon nitride, although the weight loss of the hot- pressed material was less than that of the sintered one as reported by Tressler et al. [8.45] in a previous study. Their strength data are shown in Fig. 8.4. The exposure to a eutectic mixture of NaCl and Na 2 SO 4 was more severe than to molten NaCl alone for the hot-pressed material, whereas for the reaction sintered material the effect was about the same. The differences between these two materials were attributed to the diffusion of contaminants along grain boundaries in the hot- pressed material and penetration of contaminants into pores of the reaction sintered material. This was based upon the observation that the grain boundaries of the hot-pressed material were more severely affected than those of the reaction sintered material, which did not contain an oxide grain boundary phase. The lowered fracture strengths resulted from an increase in the critical flaw size and a decrease in the critical stress intensity factor. The slight increase in fracture strengths at 1200°C was a result of a slight increase in the critical stress intensity factor. The NaCl/Na 2 SO 4 eutectic mixture, being more oxidizing than the NaCl melt, caused a greater increase in the critical flaw size. In the application of ceramics to turbine engines, the static fatigue life is of prime importance. Compared to the other types of mechanical testing in corrosive environments, little work has been reported on the long time exposure effects to Copyright © 2004 by Marcel Dekker, Inc. Properties and Corrosion 357 simulated gas turbine rig, where the corrosive environment was continued throughout the 1000°C/40 hr of the test. Room- temperature MOR fracture origins were located at pits in 17 of 22 samples. Pit formation was attributed to gas evolution during the oxidation of the silicon nitride and subsequent reaction of the silica with sodium sulfate-forming a low viscosity sodium silicate liquid. Fracture stresses were on the order of 300 MPa after exposure. Boron- and carbon-doped injected molded sintered α-SiC sprayed with thin films of Na 2 SO 4 and Na 2 CO 3 were exposed to several gas mixtures at 1000°C for 48 hr by Smialek and Jacobson [8.49]. The gas mixtures used were 0.1%SO 2 in oxygen and 0.1%CO 2 in oxygen in combination with the sulfate or carbonate thin films, respectively. The sulfate-covered sample was also exposed to pure air. Strength degradation was most severe in the sulfate/SO 2 exposure (49% loss in strength), intermediate in the sulfate/air exposure (38% loss in strength), and least severe in the carbonate/CO 2 exposure. The latter exposure caused a statistically insignificant decrease in strength when analyzed by Student’s t-test.* The primary mode of degradation was the formation of pits that varied in size and frequency depending upon the corrosion conditions. The size of the pits correlated quite well with the strength degradation (i.e., larger pits caused greater strength loss). Jacobson and Smialek [8.50] attributed this pit formation to the disruption of the silica scale by the evolution of gases and bubble formation. Zirconia-Containing Materials Although a considerable amount of scatter existed in the data of Swab and Leatherman [8.46], they concluded that Ce-TZP survived 500 hr at 1000°C in contact with Na 2 SO 4 at stress levels below 200 MPa. At stress levels greater than 250 MPa, * The application of the Student’s t-test can be found in any elementary statistics book. Copyright © 2004 by Marcel Dekker, Inc. 358 Chapter 8 failure occurred upon loading the samples. Swab and Leatherman also reported a 30% decrease in the room- temperature strength of Y-TZP after 500 hr at 1000°C in the presence of Na 2 SO 4 . This lowered strength for Y-TZP was probably a result of leaching of the yttria from the surface, which caused the transformation of the tetragonal phase to the monoclinic phase. 8.3.5 Degradation by Molten Metals The strength degradation of sintered α-silicon carbide was evaluated in both an as-received and as-ground (600 grit) condition after exposure to molten lithium by Cree and Amateau [8.51]. Transgranular fracture was exhibited for all samples when treated at temperatures below 600°C. At temperatures above 600°C, both transgranular and intergranular fracture occurred. The transgranular fracture strengths were generally greater than 200 MPa, whereas the intergranular strengths were less than 200 MPa. The low- strength intergranular failure was attributed to lithium penetration along grain boundaries beyond the depth of the uniform surface layer that formed on all samples. Grain boundary degradation was caused by the formation of Li 2 SiO 3 , from the reaction of oxidized lithium and silica. The formation of lithium silicate was accompanied by an increase in volume by as much as 25%, depending upon the temperature of exposure. The localized stresses caused by this expansion promoted intergranular crack propagation. 8.3.6 Degradation by Aqueous Solutions Bioactive Materials Bioactive ceramics include those materials that rapidly react with human tissue to form direct chemical bonds across the interface. Poor bonding across this interface and a sensitivity to stress corrosion cracking has limited the use of some Copyright © 2004 by Marcel Dekker, Inc. Properties and Corrosion 359 materials. Alumina is one material that has received a reasonable amount of study. Porous alumina has been shown to lose 35% of its strength in vivo after 12 weeks [8.52]. Seidelmann et al. [8.53] have shown that alumina loses about 15% of its strength after exposure to deionized water or blood when subjected to a constant stress. They also concluded that the service life of a hip endoprosthesis was dependent upon the density of the alumina. Ritter et al. [8.54] studied the effects of coating alumina with a bioactive glass that retarded the fatigue process. Bioactive glasses, although bonding well to bone and soft tissue, generally lack good mechanical properties. Bioactive glasses are especially sensitive to stress corrosion cracking. Barry and Nicholson [8.55] reported that a soda-lime phosphosilicate bioactive glass was unsuitable for prosthetic use at stresses above 15 MPa, thus limiting its use to tooth prostheses. This glass sustained a tensile stress of 17 MPa for only 10 years in a pH=7.4 environment. Troczynski and Nicholson [8.56] then studied the fatigue behavior of particulate and fiber-reinforced bioactive glass of the same composition. The reinforcement materials were either -325 mesh silver powder or silicon carbide whiskers. These materials were mixed with powdered glass and hot-pressed at 700°C and 30 MPa for 30 min. The composite containing the silver particulates exhibited a decreased sensitivity to stress corrosion cracking, while the composite containing the silicon carbide whiskers exhibited a sensitivity similar to that of the pure glass. Comparison of the 10-year lifetimes of the two composites indicated that the particulate-containing material survived a static stress of 22 MPa, and the whisker-containing material survived a static stress of 34 MPa. Fractography results indicated agglomerate-initiated failure for the composites as opposed to surface machining defects for the pure bioactive glass. Nitrides In the evaluation of several hot isostatically pressed silicon nitrides, Sato et al. [8.57] found that the dissolution in HCl of the sintering aids (Y 2 O 3 and Al 2 O 3 ) from the grain boundaries Copyright © 2004 by Marcel Dekker, Inc. 360 Chapter 8 decreased the three-point flexural strength. Their test variables included acid concentration, temperature, duration of dissolution, and crystallinity of the grain boundary phase. In general, the flexural strength decreased with increasing dissolution of Y 3+ and Al 3+ cations. Strengths were decreased by at least 50% after being exposed to 1 M HCl solution for 240 hr at 70°C. As expected, the grain boundary phase, having the highest degree of crystallinity, exhibited the highest strength (i.e., it is easier to leach cations from a glass than from a crystal). A control composition containing no sintering aids exhibited little, if any, strength degradation after the HCl treatment, although the strengths were considerably below those materials containing sintering aids (initially 240 vs. 600 MPa). Glassy Materials In their investigation of silica optical fibers, Dabbs and Lawn [8.58] presented data that questioned the acceptance of the Griffith flaw concept, which assumed that the flaws were exclusively cracklike and were free of preexisting influences. The real problem lies in predicting fatigue parameters for ultra- small flaws from macroscopic crack velocity data. Abrupt changes in lifetime characteristics can occur as a result of evolution of flaws long after their inception. To conduct experiments with well-defined flaws, many investigators are now using microindentation techniques. It has been reported by Lawn and Evans [8.59] that the formation of radial cracks from indentations is dependent upon the applied load. There exists a threshold load below which no radial cracks are generated; however, radial cracks may spontaneously form at the corners of subthreshold indentations long after the initial indent has been implanted if the surface is exposed to water [8.60]. Dabbs and Lawn reported data for silica optical fibers showing an abrupt increase in strength under low load conditions below the threshold for formation of radial cracks. They attributed this behavior to a transition from crack propagation-controlled failure to one of crack initiation- controlled failure. Although the subthreshold indents had no Copyright © 2004 by Marcel Dekker, Inc. [...]... influence of solubility on the reliability of optical fiber Optical Materials Reliability and Testing, SPIE 1992, 1791, 52–60 Copyright © 2004 by Marcel Dekker, Inc Properties and Corrosion 371 8.64 Matthewson, M.J.; Yuce, H.H Kinetics of degradation during aging and fatigue of fused silica optical fiber In Fiber Optics Materials and Components; Proc SPIE 1994, 2290, 204– 210 8.65 Iler, R.K The Chemistry of. .. Spear, K.E Silicon carbide materials in metallurgical heat-treatment environments Am Ceram Soc Bull 1992, 71 (11), 1683–1690 8.41 Li, C-T.; Langley, N.R Development of a fractographic method for the study of high-temperature failure of ceramic Copyright © 2004 by Marcel Dekker, Inc Properties and Corrosion 369 fibers In Advances in Ceramics, Fractography of Glasses and Ceramics; Varner, J.R Frechette,... been the use of fusion cast aluminazirconia-silica cruciform products These are in the shape of a cross and are stacked in interlocking columns This represents not only a change in chemistry, but also a change in the shape of the product, both of which lead to better overall performance A part of the concept of improvement through chemistry changes is that of improving resistance to corrosion of the bonding... Summary of results of the effects of environments on mechanical behavior of high-performance ceramics Ceram Eng Sci Proc 1991, 12 (9,10), 1886–1 913 8.27 Hench, L.L.; Ohuchi, F.; Freiman, S.W.; Wu, C.Cm.; McKinney, K.R Infrared reflection analysis of Si 3 N 4 oxidation Ceramic Engineering and Science Proceedings, 1980; Vol 1 (7–8A), 318–330 8.28 Easler, T.E.; Bradt, R.C.; Tressler, R.E Effects of oxidation... Vol 22, 442 pp Ceramic Transactions, Fractography of Glasses and Ceramics II, Frechette, V.D., Varner, J.R., Eds.; Am Ceram Soc., Westerville, OH, 1991; Vol 17, 548 pp Fracture in Ceramic Materials, Evans A.G., Ed.; Noyes Publications, Park Ridge, NJ, 1984, 420 pp 8.5 EXERCISES, QUESTIONS, AND PROBLEMS 1 Describe stress corrosion cracking and the consequences that relate to engineering materials 2 Describe... chemistry was changed to take advantage of the formation of multiple phases and the effects of surface energy upon penetrating liquids See Chap 2, Sec 2.5.3 for a discussion of the surface energy effects of multiphase systems The direct bonded basic refractory contains magnesia and spinel crystalline phases along with a grain boundary phase that is partially amorphous and partially small spinel crystals At... Simpson, W.A Influence of diesel engine combustion on the rupture strength of partially stabilized zirconia In Proceedings of the 3rd International Symposium on Ceramic Materials and Components for Engines; Tennery, V.J Ed.; Am Ceram Soc.: Westerville, OH, 1989; 549–558 8.43 Butt, D.P.; Mecholsky, J.J Effects of sodium silicate exposure at high temperature on sintered α-silicon carbide and siliconized silicon... 127 130 8.56 Troczynski, T.B.; Nicholson, P.S Stress corrosion cracking of bioactive glass composites J Am Ceram Soc 1990, 73 (1), 164–166 8.57 Sato, T.; Tokunaga, Y.; Endo, T.; Shimada, M.; Komeya, K.; Komatsu, M.; Kameda, T Corrosion of silicon nitride ceramics in aqueous hydrogen chloride solutions J Am Ceram Soc 1988, 71 (12), 1074–1079 8.58 Dabbs, T.P.; Lawn, B.R Strength and fatigue properties of. .. Chapter 8 8.31 Wang, L.; He, C.; Wu, J.G Oxidation of sintered silicon nitride materials In Proceedings of the 3rd International Symposium on Ceramic Materials and Components for Engines; Tennery, V.J., Ed.; Am Ceram Soc.: Westerville, OH, 1989; 604–611 8.32 Lange, F.F.; Davis, B.I Development of surface stresses during the oxidation of several Si3N4/CeO2 materials J Am Ceram Soc 1979, 62 (11–12), 629–630... INTRODUCTION The control of the chemical reactivity of ceramics with their environment is one of the most important problems facing the ceramics industry today Through the study of corrosion phenomena, one can learn best how to provide the control of the chemical reactivity that will provide a maximum service life expectancy at a minimum cost Most methods used to minimize corrosion have generally been . high-temperature failure of ceramic Copyright © 2004 by Marcel Dekker, Inc. Properties and Corrosion 369 fibers. In Advances in Ceramics, Fractography of Glasses and Ceramics; Varner, J.R. Frechette,. in Ceramics, Fractography of Glasses and Ceramics, Varner, J.R., Frechette, V.D., Eds.; Am. Ceram. Soc., Westerville, OH, 1988; Vol. 22, 442 pp. Ceramic Transactions, Fractography of Glasses and. exhibited oxidation of the silicon matrix and attack of the large SiC grains. In a study of the effects of molten salt upon the mechanical properties of silicon nitride, Bourne and Tressler [8.44]